Aluminum alloy containing copper and zinc

ABSTRACT

Disclosed is an elementally mixed Al—Cu—Zn base powder blend for a sintered Al base alloy. The powder blend includes more than 5.6 wt % and less than 9 wt % Cu, 1˜5 wt % Zn, and a balance Al. With the powder blend, an article of a sintered Al base alloy having higher wear resistance as well as higher tensile strength can be fabricated.

CROSS-REFERENCES TO RELATED APPLICATIONS

The present application claims priority to and the benefit of Korean Patent Application No. 10-2006-0010896, filed Feb. 4, 2006, the disclosure of which is incorporated by reference in its entirety.

BACKGROUND

1. Field

The present invention relates to an aluminum alloy, and more particularly, to a sintered aluminum base alloy.

2. Discussion of Related Technology

Recently, Al-base alloys have been actively replacing various ferrous components in automobiles to reduce the weight and improve the performance. (See The International Journal of Powder Metallurgy, 36, 2002, p 41 by F. V. Beaumont, and p 45 by C. Lall et al.) Most of aluminum alloys can be easily processed to final shapes via casting, forging, forming or machining, and also heat-treated to improve desired final properties. Many automotive components, such as space frames, an engine blocks, wheel frames, housings, etc, are currently utilizing aluminum castings or forgings. For ferrous or copper-base components in automobiles, powder metallurgy products are widely used because of their easy net-shaping capability, cost competitiveness and acceptable properties. Aluminum base powder products, however, have found very limited applications in automobiles despite many potential merits. (See The International Journal of Powder Metallurgy, 36, p 51, 2002 by W. H. Hunt) One of the main obstacles limiting the use of aluminum sintered products is the poor sintering behavior of aluminum-base powders which cause relatively low sintered properties. (See Nature, 181, p 833, 1958 by R. F. Smart et al., and Journal of Metals, 37, p 27, 1985 by Y. Kim et al.)

In general, aluminum-base powder mixtures for sintered products are prepared by mixing air atomized aluminum powder with secondary alloying powders such as copper, magnesium, silicon, zinc or others. The air atomized aluminum powder, though least expensive, contains a relatively thick aluminum oxide layer around each particle surface due to the oxidation during the atomization. Since the aluminum oxide is very stable and the sintering temperature for aluminum is relatively low, it is very difficult to reduce the oxide layer during sintering at around 600° C. The oxide layers between particles in a compact block the inter-diffusion between aluminum particles, thus severely limiting the sintering process. Compaction of a powder mixture at a relatively high pressure in order to break the surface oxide layers is known to improve the sintering to some extent.

However, forming a liquid phase during the sintering is considered as the most effective way of improving the sintering of elementally mixed aluminum powder compacts. (See “Properties and design guidelines for aluminum parts” in Proceeding 2000 International Conference on P/M Aluminum & Light Alloys for Automotive Applications, pp 51-58 by Antonio Romero, Acta Materialia, 35, p 589, 1996 by R. N. Lumley et al., and Material Chemistry and Physics, 67, p 85, 2001 by G. B. Schaffer et al.) Such liquid phases can be formed easily during the heating period of sintering by eutectic reaction between the aluminum powder and additive elemental powders such as copper, zinc, magnesium, silicon, etc. When a powder compact is heated to sintering temperature, localized alloying occurs at the contacts between aluminum and the additive powders to form a liquid phase during the heating cycle or sintering. The liquid phase, if persistently present during the sintering, can spread through particle boundaries, pulling particles and filling pores, so that higher sintered density and better bonding between particles can be obtained. On the other hand, if the liquid phase is transient, that is, formed at a relatively low temperature but disappeared by solutionizing in the matrix during heating or in the early stage of sintering, the dissolved liquid phase leaves pores behind and little persistent liquid phase in the matrix to assist particle bonding during the sintering. Therefore, controlling the characteristics and the amount of the liquid phase are most critical to improve the sintering of mixed elemental aluminum powders.

Most of commercially available mixed elemental aluminum powder blends, however, have the compositions based on those of wrought aluminum-base alloy systems, i.e., 201AB of Alcoa Inc. and AMB2712 of Ampal Inc. from AA2014, 601AB of Alcoa Inc. and AMB6712 of Ampal Inc. from AA6061, and AMB 7775 of Ampal Inc. from AA7075, rather than being optimized for sintering. The sintered properties obtainable from these powder mixtures are relatively low when compared to those of the wrought counterparts. For example, sintered Alcoa 201AB at 95% theoretical density was reported to have the tensile strength of ˜330 MPa after T6 condition, which is about only 70% of that of AA2014, despite the only 5% of porosity level. (See above “Properties and design guidelines for aluminum parts” in Proceeding 2000 International Conference on P/M Aluminum & Light Alloys for Automotive Applications, pp 51-58 by Antonio Romero) AMB7775 was reported to have the highest tensile strength over 400 MPa, but exhibits a relatively poor wear resistance. AMB7775 contains 6-8 wt % Zn, 3-5 wt % Mg and 1-3 wt % Cu with a small amount of Si and Pb, and also forms a liquid phase during the sintering. The liquid phase, however, is mainly transient, that is, disappeared by dissolving in the aluminum matrix during the sintering because of relatively large solubility of these elements in aluminum at or below the sintering temperature. The blend thus can be well strengthened by solid solution hardening and also by precipitation hardening, but dimensional control could be rather difficult because of the pores left by the solutionized liquid phase.

U.S. Pat. No.5,902,943 describes about Al—Zn-Mg—Cu base mixed elemental aluminum alloy powder blends and sintered aluminum alloys. The blends described in the patent are basically 7xxx type alloys and contain a relatively large amount of Zn as the principal alloying addition and other elements to enhance the sintering. Depending on the relative quantities of the alloying additions and heat treatment conditions, the blends can exhibit tensile strengths over 400 MPa but require a precise control of processing variables. The patent does not describe any results on the wear properties of the blends. Wear is also a big concern for most of potential aluminum powder metallurgy (PM) parts which will replace steel parts. Aluminum alloys generally have a relatively poor wear resistance than steels and thus require methods to improve the wear resistance. Hypereutectic Al—Si alloys which contain over 20% Si possess excellent wear resistance when they are produced in a pre-alloyed powder form and consolidated to full density at elevated temperatures. The pre-alloyed powders, however, are not well suited for the press and sinter processing, and thus are not very economical. (See above Material Chemistry and Physics, 67, p 85, 2001 by G. B. Schaffer et al.) Aluminum alloy composites reinforced with hard ceramic particles are alternatives but they are generally not well sinter-able and thus possess very poor strength and ductility in as-sintered condition.

Therefore, new mixed elemental aluminum base powder blends which can provide a better combination of strength and wear resistance are needed for wider range of applications than has been possible with existing blends.

The discussion in this section is to provide general background information, and does not constitute an admission of prior art.

SUMMARY

One aspect of the invention provides an aluminum alloy comprising Al, Cu and Zn, wherein a portion of the alloy comprises: Cu in an amount over 5.6 wt % and less than about 9 wt % with reference to the weight of the portion; and Zn in an amount from about 1 wt % to about 5 wt % with reference to the weight of the portion.

In the foregoing alloy, the portion may comprise Al-containing grains and an intergrain material disposed between and interconnecting neighboring grains, wherein a substantial amount of Cu may be present in the intergrain material, and wherein a substantial amount of Zn is present in central areas of the Al-containing grains. In one embodiment, the substantial amount of Cu present in the intergrain material is about 2.5% to about 95% of the total amount of Cu. In certain embodiments, Cu present in the intergrain material may be about 2.5, 5, 7.5, 10, 15, 20, 25, 30, 35, 40, 50, 65, 80 or 95% of its total amount in the alloy portion. In some embodiments, Cu present in the intergrain material may be within a range defined by two of the foregoing amounts. In one embodiment, the substantial amount of Zn present in the Al-containing grain is about 5% to 100% of the total amount of Zn. In certain embodiments, Zn present in the Al-containing grain may be about 5, 20, 35, 50, 60, 70, 80, 85, 90, 95, 97, 98, 99, 99.5 or 100% of its total amount in the alloy portion. In some embodiments, Zn present in the Al-containing grain may be within a range defined by two of the foregoing amounts. Substantially the entire amount of Zn may be present in the Al-containing grains.

Still in the foregoing alloy, a substantial amount of Cu present in the intergrain material is in the form of CuAl₂. Substantially the entire amount of Cu present in the intergrain material may be in the form of CuAl₂. In one embodiment, the substantial amount of Cu which is present in the intergrain material in the form of CuAl₂ is about 5% to 100% of the total amount of Cu present in the intergrain material. In certain embodiments, Cu present in the intergrain material in the form of CuAl₂ may be about 5, 20, 35, 50, 60, 70, 80, 85, 90, 95, 97, 98, 99, 99.5 or 100% of its total amount in the intergrain material. In some embodiments, Cu present in the intergrain material in the form of CuAl₂ may be within a range defined by two of the foregoing amounts. The intergrain material may comprise a portion which comprises Cu in an amount from about 20 wt % to about 60 wt % with reference to the weight of the portion of the intergrain material.

Further in the foregoing alloy, the portion of the alloy may comprise Cu in an amount from about 6 wt % to about 8 wt % with reference to the weight of the portion. The portion of the alloy may contain Sn in an amount from about 0.01 wt % to about 0.05 wt % with reference to the weight of the portion. The portion of the alloy may contain Mg in an amount less than about 0.03 wt % with reference to the weight of the portion. The alloy may be produced by a method, which comprises: providing a powder mixture comprising Al, Cu and Zn; and heating the powder mixture to a temperature sufficient to melt at least part of the powder mixture.

Another aspect of the invention provides a method of making an aluminum alloy, comprising: providing a powder mixture comprising Al, Cu and Zn, wherein Cu is in an amount over 5.6 wt % and less than about 9 wt % with reference to the weight of the powder mixture, and wherein Zn is in an amount from about 1 wt % to about 5 wt % with reference to the weight of the powder mixture; and heating the powder mixture to a temperature sufficient to melt at least part of the powder mixture.

In the foregoing method, the powder mixture may comprise Al-containing particles, and wherein a substantial amount of Zn may be dissolved into at least part of the Al-containing particles. Upon heating at least part of the powder mixture is melt to form a liquefied state, and wherein a substantial amount of Cu may be present in the liquefied state. The method may further comprise cooling the heated powder mixture thereby forming an alloy comprising a portion which comprises Al-containing grains and an intergrain material disposed between and interconnecting neighboring grains, wherein a substantial amount of Cu is present in the intergrain material.

Yet another aspect of the invention provides an aluminum alloy produced by the foregoing method, wherein a portion of the alloy comprises Al-containing grains and an intergrain material disposed between and interconnecting neighboring grains, wherein a substantial amount of Cu is present in the intergrain material, and wherein a substantial amount of Zn is present in central areas of the Al-containing grains.

A further aspect of the invention provides a powder blend for use in making an aluminum alloy, the powder blend comprising Al, Cu and Zn, wherein the powder blend comprises: Cu in an amount over 5.6 wt % and less than about 9 wt % with reference to the weight of the powder blend; and Zn in an amount from about 1 wt % to about 5 wt % with reference to the weight of the powder blend.

In the foregoing powder blend, wherein the powder blend may comprise Cu in an amount from about 6 wt % to about 8 wt % with reference to the weight of the powder blend. The powder blend may comprise Sn in an amount from about 0.01 wt % to about 0.05 wt % with reference to the weight of the powder blend. The powder blend may comprises Mg in an amount less than about 0.03 wt % with reference to the weight of the powder blend.

One or more embodiments of the present invention provide an elementally mixed Al—Cu—Zn base powder blend, a method of fabricating an article of a sintered alloy using the powder blend, and an article fabricated using the powder blend.

An aspect of the present invention provides a powder blend comprising more than 5.6 wt % Cu added to a balance Al to form a mixed powder blend. Thus, considerable amount of a liquid phase, which is formed over the eutectic temperature of Al—Cu, i.e. 548° C., persistently presents at a sintering temperature (about 600° C.), though a portion of the liquid phase is solutionized into a matrix. The persistent liquid phase fills boundaries and pores between powders as well as accelerates the sintering of the solid powders. Thus densification of the sintered alloy is improved.

Maximal solid solubility of Cu in an Al matrix is about 5.5˜5.6 wt % at the eutectic temperature 548° C. Thus, when more than 5.6 wt % Cu powder is added to Al powder, the persistent liquid phase always presents during the sintering, without relation to a temperature elevation rate, etc. The liquid phase which persistently presents during the sintering is solidified into a mixed phase of α-Al and CuAl₂ (θ) phase, as the liquid phase is cooled to the room temperature, and the solidified liquid phase incorporates Al and 35 wt % Cu. (See Journal of Materials Science, 40, p 441, 2005 by Sang Chul et al.) The CuAl₂ phase has vickers hardness (HV) of 980 higher than that of α-Al. (See Intermetallics, 7 p 1001, 1999 by D. Moreno et al.) The CuAl₂ phase contributes to improve both of the strength and the wear resistance of the sintered alloy. However, increment of the amount of the CuAl₂ phase may result in decrease of ductility of the sintered alloy. Thus, the amount of Cu may be decided by way of considering the strength, ductility, productivity of a sintered article, required shape of the article, required dimensions and tolerance of the article and deformation of a product shape during sintering. It is preferable to limit the amount of Cu to less than 9 wt %.

Meanwhile, since the solid solubility of Cu at a temperature of about 600° C. for sintering Al base powder blends is less than 3 wt %, the effect of solid solution strengthening is limited. Thus, Zn powder which has a higher solid solubility is added to improve the solid solution strengthening effect. Zn reacts with Al during a heating period of sintering, and forms a eutectic liquid phase over 382° C., and then is all solutionized into the Al matrix to enhance the strength of the matrix. In addition, Zn together with Cu can improve a hardening effect resulting from an age-hardening treatment after a solution-treatment, thus enhancing the strength of a sintered alloy. Increment of addition of Zn increases the strength of a sintered alloy and a thermal treated sintered alloy. However, an excessive transient liquid phase formed due to the surplus increment of Zn makes it difficult to maintain the shape of a compact. Therefore, it is preferable to limit the addition of Zn to not more than 5 wt %.

An aspect of the present invention provides a method of fabricating an article of a sintered Al—Cu—Zn base alloy. The method comprises mixing more than 5.6 wt % and less than 9 wt % Cu, 1˜5 wt % Zn, and a balance Al to form a mixed powder blend. The mixed powder blend is compacted, and then the compacted powder is sintered to form a sintered alloy. As a result, an article of a sintered Al—Cu—Zn base alloy can be fabricated.

The sintered alloy may be easily re-pressed because it has low yield strength (YS). With the re-pressing, a sintered alloy having a theoretical density of over 95% can be obtained, and the increase of the density improves the strength of the article of the sintered alloy. Also, deformation occurred during the sintering can be corrected using the re-pressing. Thus, the article having precise dimensions can be fabricated. In addition, the re-pressed alloy can be heat-treated to fabricate an article of a sintered alloy having a better combination of strength and wear resistance. The heat treatment may comprise solution-treating the re-pressed alloy, and heat-treating the solution-treated alloy for an age-hardening. Usually, the solution-treated alloy is water-quenched before the heat treating for an age-hardening. With the solution treatment, Cu in the liquid phase is solutionized into an Al matrix, and then CuAl_(2-x) (θ″, θ′, or θ) phases are precipitated by the age-hardening treatment. As a result, the strength and hardness of the sintered alloy are improved.

An aspect of the present invention provides an article of a sintered Al—Cu—Zn base alloy. The article includes more than 5.6 wt % and less than 9 wt % Cu, 1˜5 wt % Zn, and a balance Al. The article of the sintered Al—Cu—Zn base alloy may be fabricated using the fabricating method described above. Meanwhile, the article of the sintered Al—Cu—Zn base alloy comprises an Al matrix (α-Al) and a CuAl₂ phase. The CuAl₂ phase may present at boundaries of α-Al and/or in the Al matrix. The CuAl₂ phase serves as a reinforcing phase and improves both of the strength and the wear resistance of the article of the sintered Al—Cu—Zn base alloy. In several embodiments of the present invention, additive powders may be incorporated in the Al—Cu—Zn base powder blends. Mg added to the Al—Cu—Zn powder blends decreases the strength and ductility of its sintered alloy, whereas a small amount of Sn added to the Al—Cu—Zn powder blends increases the ductility of its sintered alloy. Preferably, 0.010.05 wt % Sn may be added to the powder blends, and less than 0.01 wt % Mg may be incorporated in the powder blends.

BRIEF DESCRIPTION OF THE DRAWINGS

The above features and advantages of the present invention will become more apparent by describing the embodiment thereof with reference to the accompanying drawings, in which:

FIG. 1 is a schematic diagram depicting a method of measuring transverse rupture strength (TRS) and an amount of deflection at rupture of a sintered sample.

FIG. 2 is a graph for describing TRSs and amounts of deflection of sintered samples of Al base alloys incorporating different Cu contents.

FIG. 3 shows SEM images of sintered Al—Cu base alloys incorporating different Cu contents according to certain embodiments and a sintered Alcoa 201AB.

FIG. 4 is a graph depicting TRSs and amounts of deflection of sintered samples of Al-6 wt % Cu powder mixtures with sintering temperatures.

FIG. 5 is a graph depicting TRSs and amounts of deflection of sintered samples with different Zn contents in Al—Cu—Zn powder blends.

FIG. 6 shows optical images of sintered samples with different Zn contents in Al—Cu—Zn powder blends.

FIG. 7 shows a SEM image for describing a solidified liquid phase (A) and a matrix (B) of a sintered Al—Cu—Zn base alloy.

FIG. 8 is a graph depicting TRSs and amounts of deflection of sintered samples with different Mg contents in Al—Cu powder blends.

FIG. 9 shows optical images of sintered samples with different Mg contents in Al—Cu powder blends.

FIG. 10 is a graph depicting TRSs and amounts of deflection of sintered samples with different Mg contents in Al—Cu—Zn powder blends.

FIG. 11 is a graph depicting TRSs and amounts of deflection of sintered samples with different Sn contents in Al—Cu—Zn powder blends.

FIG. 12 is a graph depicting age-hardening behaviors of sintered Al—Cu—Zn base alloys with aging temperature and time.

FIG. 13 shows a SEM image of a heat-treated sintered Al—Cu—Zn base alloy

FIG. 14 shows XRD results of sintered alloys of an Al—Cu—Zn base powder blend and a commercially available powder blend.

DETAILED DESCRIPTION OF EMBODIMENTS

Hereinafter, embodiments of the invention will be described in detail with reference to the accompanying drawings.

Table 1 lists typical compositions and tensile properties of commercially available aluminum base mixed elemental powder blends. 2xxx blends contain Cu as a main additive element, however, Cu is less than 5 wt %, and 7xxx blends contain Zn as a main additive element. All of the commercially available aluminum base powder blends show a transient liquid phase sintering behavior, that is, the liquid phase formed during the sintering is almost solutionized or absorbed into a matrix. Among these powder blends, 7xxx series possess the highest tensile strength at both as-sintered and heat-treated conditions. Wherein, T1 indicates the as-sintered condition, and T6 indicates the heat-treated condition, that is, aged to peak hardness condition. TABLE 1 Tensile properties Composition(wt %) Elongation wrought Cu Mg Si Zn Al YS(MPa) TS(MPa) (%) equivalent 6xxx: 0.2 1.0 0.5 — Bal. 135(T1) 5 6061 Al—Mg—Si 240(T6) 3 2xxx: 4.4 0.5 0.7 — Bal. 180(T1) 205(T1) 5 2014 Al—Cu—Mg 327(T6) 330(T6) 1 7xxx: 1.5 2.5 — 5.5 Bal. 230(T1) 270(T1) 4 7075 Al—Zn—Mg—Cu 370(T6) 413(T6) 2

FIG. 1 shows how to measure transverse rupture strength (TRS) and an amount of deflection at rupture of a sintered sample in accordance with ASTM B312. The TRS and the amount of deflection are utilized as the measure of mechanical strength and ductility of the sintered sample, respectively. Thickness, length and width of sample for this test was 6.35 mm, 31.8 mm and 12.7 mm, respectively.

FIG. 2 is a graph for describing TRSs and amounts of deflection of sintered samples of Al base alloys incorporating different Cu contents. Each sample is provided by sintering a mixed powder blend of Al and Cu powders at 600° C. for 1 hour in a dry nitrogen atmosphere, and then slowly cooling the sintered alloy to room temperature. For comparison, two commercial powder blends, 201AB and 601AB of Alcoa Inc., were processed via same procedure and their TRSs and deflections were also measured. Referring to FIG. 2, it is noticeable that with increasing the Cu content, TRS increases but ductility decreases. In addition, all the samples incorporating 6, 8 and 10 wt % Cu showed much higher strength than the samples of the commercial powder blends, and showed ductility almost similar to the samples of the commercial powder blends.

FIG. 3 shows SEM images of sintered alloys incorporating different Cu amount and a sintered Alcoa 201AB, where (a) is for Al-6 wt % Cu, (b) is for Al-8 wt % Cu, (c) is for Al-10 wt % Cu, and (d) is for Alcoa 201AB. Referring to FIG. 3, a liquid phase is observed in the samples of (a), (b) and (c). The solidified liquid phase fills boundaries and pores between powders, and increases with increasing the Cu content. The solidified liquid phase can be an evidence of a persistent liquid phase during the sintering. However, (d) for the sintered Aloca 201AB shows some coarse pores with no evidence of the persistent liquid phase despite of about 4.5 wt % of Cu content. In certain embodiments, the amount of Cu may be about 5.6, 5.7, 5.8, 6, 6.2, 6.5, 7, 7.5, 8, 8.5, 9 or 9.5 wt % of the sintered alloy. In some embodiments, the amount of Cu may be within a range defined by two of the foregoing amounts.

FIG. 4 is a graph depicting TRSs and amounts of deflection of sintered samples of Al-6 wt %Cu (Al-6Cu) powder mixtures with sintering temperatures. Each power blend is sintered for 1 hour in a dry nitrogen atmosphere, and then slowly cooled to room temperature. Referring to FIG. 4, the strength and the amount of deflection at rupture increased with the increase of the sintering temperature. And, for both of the strength and deflection, marked differences were observed between 590° C. and 600° C. This is attributed to the significant improvement of sintering due to an extensive liquid phase formation at temperatures of or above 600° C.

FIG. 5 is a graph depicting TRSs and amounts of deflection of sintered samples with different Zn contents in Al—Cu—Zn powder blends. All samples contained 6 wt % Cu, and were sintered at 625° C. for 1 hour in a dry nitrogen atmosphere. Referring to FIG. 5, even 1 wt % of Zn addition showed a quite significant effect for increasing the strength. The strength also increased with the increase of Zn, while the ductility decreased slightly. This is a result of solid solution strengthening effect due to the solid solution of Zn in Al matrix or grains.

FIG. 6 shows optical images of sintered samples with different Zn contents in Al—Cu—Zn powder blends. Where, (a) is for Al-6Cu-1Zn, (b) is for Al-6Cu-3Zn, and (c) is for Al-6Cu-5Zn. All samples were sintered at 625° C. for 1 hour in a dry nitrogen atmosphere. Referring to FIG. 6, with increasing the Zn content, the amount of pore decreased and grains became larger. In certain embodiments, the amount of Zn may be about 0.5, 1, 2, 3, 4, 4.5, 5 or 5.5 wt % of the sintered alloy. In some embodiments, the amount of Zn may be within a range defined by two of the foregoing amounts.

FIG. 7 shows a SEM image for describing an intergrain material or solidified liquid phase (A) and a Al matrix or grains (B) of a sintered Al—Cu—Zn base alloy. Compositional variations of liquid phase (A) and the Al matrix (B) were analyzed using an EDAX, and the results are listed in Table 2. The sintered alloy was provided by sintering an Al-6Cu-3Zn powder blend at 625° C. for 1 hour in a dry nitrogen atmosphere. TABLE 2 A B Element wt % at % wt % at % O 1.05 2.36 1.3 2.28 Al 54 72.14 90.46 94.12 Cu 44.95 25.5 4.49 1.98 Zn 0 0 3.75 1.61 Totals 100 100 100 100

Referring to Table 2, the solidified liquid phase (A) includes Al and Cu and does not incorporate Zn. The solidified liquid phase was identified to consist of α-Al and CuAl₂ (θ) phase by XRD analysis of FIG. 14, and their relative fractions were calculated with lever rule as about 15 wt % and about 85 wt %, respectively. All of the Zn along with considerable amounts of Cu was solutionized in the matrix, thus improving a solid solution strengthening effect. Therefore, the observed microstructure of FIG. 7 can be regarded as a composite material where Al—Cu—Zn solid solutionized matrix is strengthened by the reinforcing CuAl₂ phase which contributes to increased strength and wear resistance of the alloy. (B) shows a contents of Al matrix or grains.

FIG. 8 is a graph depicting TRSs and amounts of deflection of sintered samples with different Mg contents in Al-6Cu powder blends. And, FIG. 9 shows optical images of sintered samples with different Mg contents in Al-6Cu powder blends. Where, (a) is for Al-6Cu-0.1Mg, (b) is for Al-6Cu-0.3Mg, and (c) is for Al-6Cu-0.5Mg. Referring to FIG. 8, even 0.1 wt % Mg in Al—Cu powder blends lowered the TRS and ductility. The TRSs and ductility decreased with increase of the Mg contents in Al-6Cu. Referring to FIG. 9, it is noticeable that the addition of Mg to the Al—Cu powder blends altered the sintered microstructures significantly to have many isolated pores. Increase of Mg in the powder blend resulted in more porosity and coarse pores.

FIG. 10 is a graph depicting TRSs and amounts of deflection of sintered samples with different Mg contents in Al-6Cu-3Zn powder blends. In Al—Cu—Zn systems, even 0.03 wt % Mg lowered the TRS and ductility of a sintered alloy. Thus, addition of Mg in Al—Cu—Zn powder blends is preferable to be limited below 0.03 wt %.

FIG. 11 is a graph depicting TRSs and amounts of deflection of sintered samples with different Sn contents in Al-6Cu-3Zn powder blends. In Al—Cu—Zn systems, 0.01 wt % Sn increased ductility of a sintered alloy, however, 0.05 wt % Sn decreased the TRS and ductility. Thus, 0.01˜0.05 wt % Sn may be utilized as a minor addition in the Al—Cu—Zn powder blends to control ductility of sintered alloys.

FIG. 12 is a graph depicting age-hardening behaviors of sintered Al—Cu—Zn base alloys with aging temperature and time. Sintered samples of Al-6Cu-3Zn were used, and the sintered samples were re-pressed and solution-treated at 540° C. and water-quenched, and then aging-treated at various temperatures and times. Vickers hardness (HV) and Rockwell hardness (HRB) were together shown.

Referring to FIG. 12, with increase of the aging time after the solution treatment, the hardness of the samples increased. Also, with increase of the aging temperature, the hardness of the samples increased much more. This implies that the sintered Al—Cu—Zn base alloy according to embodiments of the present invention can be effectively strengthened by heat treatments such as the solution treatment and age-hardening treatment.

FIG. 13 shows a SEM image of a heat-treated sintered Al—Cu—Zn base alloy. A sintered sample of Al-6Cu-3Zn was used, and the sintered sample was re-pressed and heat-treated. Compositional variations of liquid phase (A) and the Al matrix (B) were analyzed using an EDAX, as described above referring to FIG. 7, and the results are listed in Table 3. TABLE 3 A B Element wt % at % wt % at % Al 65.97 82.03 89.99 95.53 Cu 34.03 17.97 7.01 3.16 Zn 0 0 3 1.31 Totals 100 100 100 100

In certain embodiments, the Cu content of the Al matrix can be slightly increased after the heat treatment, but Zn content was almost unchanged. Referring to Table 3, the solidified liquid phase was found to contain about 34 wt % Cu. Therefore, the Al matrix is strengthened by fine precipitates after aging and also by the presence of a Cu-rich hard phase in. The Cu-rich hard phase consists of about 40 wt % of a-Al and 60 wt % of CuAl₂ (θ) phase. The CuAl₂ phase functions as a reinforcing phase, thus improving the strength and wear resistance of the sintered alloy.

FIG. 14 shows XRD graphs of a sintered alloys after sintering, solution-treatment and water-quenching, and then artificial aging. Where, (a) shows the XRD graphs of commercial 7xxx and (b) shows the XRD graphs of Al-6Cu-3Zn according to one embodiment of the present invention. XRD datum were obtained after sintering, after solution treatment, and after aging treatment, respectively. The XRD data after solution treatment was obtained from a sample which was solution-treated at 540° C. for 1 hour and then water-quenched. The XRD data after aging treatment was obtained from a sample which was aging-treated after the solution treatment and water-quenching.

Referring to FIG. 14, for the Al-6Cu-3Zn, main constituents of the as-sintered sample are α-Al and CuAl₂ (θ) phase. The CuAl₂ (θ) phase disappeared after the solution treatment. This is attributed to super-saturation of Cu atoms in the Al matrix after the solution-treatment at 540° C. and subsequent water-quenching. During natural or artificial aging, CuAl_(2-x) (θ′, or θ″) and CuAl₂ (θ) phases precipitate in the Al matrix (B) and the solidified liquid phase (A), respectively, effectively improving the strength and wear resistance. In the meantime, for the commercial 7xxx, the CuAl₂ (θ) phase did not appear after the aging treatment.

Table 4 lists hardness, transverse rupture and tensile properties of Al-6Cu-3Zn and Al-6Cu-5Zn mixed elemental powder alloys compacted to 90% theoretical density (T.D.) at room temperature using double action die, sintered at 610° C. for 1 hour under flowing N2 atmosphere to about 96% T.D., further re-pressed to 98% T.D., and finally heat-treated for age-hardening. All samples were maintained at 540° C. for 1 hour and then water-quenched before aging. TABLE 4 Transverse Rupture Properties Tensile Properties Deflection Yield Tensile Alloy Heat Hardness Strength at Rupture Strength Strength Elongation System Treatment (HV) (MPa) (mm) (MPa) (MPa) (%) Al—6Cu—3Zn As-sintered 64 371 1.9 60 197 10.9 150° C./19 hrs 152 673 0.69 287 410 6.9 170° C./13 hrs 152 667 0.54 317 419 5.1 Al—6Cu—5Zn As-sintered 65 374 1.7 75 204 10.2 150° C./22 hrs 152 636 0.64 314 396 4.0 170° C./10 hrs 154 642 0.55 320 399 4.1

Referring to Table 4, as-sintered samples usually show very low hardness and yield strength (YS) values and significant ductility, facilitating further plastic working process such as re-pressing and other cold working processes. Heat-treated samples usually show significant increases in hardness and strength with accompanying decrease in ductility. This is caused by precipitation of fine θ′ phase in the α-Al matrix. Strength and ductility slightly decreased by increasing Zn contents from 3 to 5 wt %. Thus, it is preferable to limit the Zn contents within 5 wt %.

Table 5 details wear resistance characteristics of various Al-base alloy systems, commercial and developed ones alike, compacted to 90% T.D., sintered for 1 hour at 610° C. under flowing N₂ atmosphere to about 96% T.D., further re-pressed to 98% T.D., and finally heat-treated for age-hardening. For age-hardening, sintered samples were solution-treated for 1 hour at 540° C. and water-quenched immediately afterward, followed by artificial ageing for 22 hours at 150° C. Wear resistance was characterized by weight loss of pins after sliding 2000 m against rotating disk at 100° C. in a commercial engine oil. The pins were pressed against the rotating disk with the force of 500 N. Both the pin and disk were made with same blend. TABLE 5 Weight Coefficient Loss of Weight Loss of Total Weight of Alloy Systems Disk Pin Loss Friction AMB7775 0.0220 0.1595 0.1815 0.6601 AMB7777- 0.0580 0.2590 0.3170 1.2515 10v/oSiC Al—4Cu—5Zn 0.1145 0.1685 0.2830 0.1289 Al—4Cu—7Zn 0.0420 0.1013 0.1433 0.1442 Al—6Cu—5Zn 0.0254 0.0325 0.0579 0.1425 Al—8Cu—5Zn 0.0365 0.0205 0.0570 0.0994

Referring to Table 5, Very high coefficients of friction were observed for commercially available 7xxx series alloys and the 7xxx alloys with reinforcing SiC particles, resulting in significant amounts of wear. Al-4Cu alloy system shows still significant amount of wear, although somewhat reduced compared to commercially available 7xxx series alloys. Al-6Cu alloy system, on the other hand, caused marked reduction in both coefficient of friction and amount of wear. This is attributed to the solidified liquid phase which acted as reinforcing phase upon solidification as illustrated in FIG. 9. Although Al-4Cu alloys show somewhat reduced coefficient of friction by the liquid phase, the contribution is regarded as minimal at best due to small amount of liquid formation in the Al-4Cu alloy. However, significant increase in wear resistance observed for Al-6Cu alloy is attributed to the ample amount of liquid phase formed during sintering.

It will be appreciated that many changes and modifications can be made to the discussed embodiments without departing from the scope of the present invention, which is defined in the following claims. 

1. An aluminum alloy comprising Al, Cu and Zn, wherein a portion of the alloy comprises: Cu in an amount over 5.6 wt % and less than about 9 wt % with reference to the weight of the portion; and Zn in an amount from about 1 wt % to about 5 wt % with reference to the weight of the portion.
 2. The alloy of claim 1, wherein the portion comprises Al-containing grains and an intergrain material disposed between and interconnecting neighboring grains, wherein a substantial amount of Cu is present in the intergrain material, and wherein a substantial amount of Zn is present in central areas of the Al-containing grains.
 3. The alloy of claim 2, wherein a substantial amount of Cu present in the intergrain material is in the form of CuAl₂.
 4. The alloy of claim 3, wherein substantially the entire amount of Cu present in the intergrain material is in the form of CuAl₂.
 5. The alloy of claim 2, wherein the intergrain material comprises a portion which comprises Cu in an amount from about 20 wt % to about 60 wt % with reference to the weight of the portion of the intergrain material.
 6. The alloy of claim 2, wherein substantially the entire amount of Zn is present in the Al-containing grains.
 7. The alloy of claim 1, wherein the portion of the alloy comprises Cu in an amount from about 6 wt % to about 8 wt % with reference to the weight of the portion.
 8. The alloy of claim 1, wherein the portion of the alloy contains Sn in an amount from about 0.01 wt % to about 0.05 wt % with reference to the weight of the portion.
 9. The alloy of claim 1, wherein the portion of the alloy contains Mg in an amount less than about 0.03 wt % with reference to the weight of the portion.
 10. The alloy of claim 1, wherein the alloy is produced by a method, which comprises: providing a powder mixture comprising Al, Cu and Zn; and heating the powder mixture to a temperature sufficient to melt at least part of the powder mixture.
 11. A method of making an aluminum alloy, comprising: providing a powder mixture comprising Al, Cu and Zn, wherein Cu is in an amount over 5.6 wt % and less than about 9 wt % with reference to the weight of the powder mixture, and wherein Zn is in an amount from about I wt % to about 5 wt % with reference to the weight of the powder mixture; and heating the powder mixture to a temperature sufficient to melt at least part of the powder mixture.
 12. The method of claim 11, wherein the powder mixture comprises Al-containing particles, and wherein a substantial amount of Zn is dissolved into at least part of the Al-containing particles.
 13. The method of claim 11, wherein upon heating at least part of the powder mixture is melt to form a liquefied state, and wherein a substantial amount of Cu is present in the liquefied state.
 14. The method of claim 11, further comprising cooling the heated powder mixture thereby forming an alloy comprising a portion which comprises Al-containing grains and an intergrain material disposed between and interconnecting neighboring grains, wherein a substantial amount of Cu is present in the intergrain material.
 15. An aluminum alloy produced by the method of claim 11, wherein a portion of the alloy comprises Al-containing grains and an intergrain material disposed between and interconnecting neighboring grains, wherein a substantial amount of Cu is present in the intergrain material, and wherein a substantial amount of Zn is present in central areas of the Al-containing grains.
 16. The alloy of claim 15, wherein a substantial amount of Cu in the intergrain material is present in the form of CuAl₂.
 17. The alloy of claim 16, wherein substantially the entire amount of Cu in the intergrain material is present in the form of CuAl₂.
 18. The alloy of claim 15, wherein a substantial amount of Zn is present in the Al-containing grains.
 19. The alloy of claim 18, wherein substantially the entire portion of Zn is present in the Al-containing grains.
 20. A powder blend for use in making an aluminum alloy, the powder blend comprising Al, Cu and Zn, wherein the powder blend comprises: Cu in an amount over 5.6 wt % and less than about 9 wt % with reference to the weight of the powder blend; and Zn in an amount from about 1 wt % to about 5 wt % with reference to the weight of the powder blend.
 21. The powder blend of claim 20, wherein the powder blend comprises Cu in an amount from about 6 wt % to about 8 wt % with reference to the weight of the powder blend.
 22. The powder blend of claim 20, wherein the powder blend comprises Sn in an amount from about 0.01 wt % to about 0.05 wt % with reference to the weight of the powder blend.
 23. The powder blend of claim 20, wherein the powder blend comprises Mg in an amount less than about 0.03 wt % with reference to the weight of the powder blend. 